Method for producing an iron-chromium alloy

ABSTRACT

The invention relates to a method for producing a component, made of an iron-chromium alloy that precipitates Laves phases and/or particles containing Fe and/or particles containing Cr and/or particles containing Si and/or carbides, by subjecting a semi-finished product made of the alloy to a thermomechanical treatment, wherein in a first step, the alloy is solution heat treated at temperatures≧the solution heat treatment temperature and is subsequently quenched in stationary protective gas or air, moving (blown) protective gas or air, or water. In a second step, a mechanical forming of the semi-finished product in a range from 0.05 to 99% is performed, and in a subsequent step, Laves phases Fe 2 (M, Si) or Fe 7 (M, Si) 6  and/or particles containing Fe and/or particles containing Cr and/or particles containing Si and/or carbides are precipitated in a specific and finely distributed manner in that the component produced from the formed semi-finished product is brought to an application temperature between 550° C. and 1000° C. by means of heating at 0.1° C./min to 1000° C./min.

The invention relates to a ferritic iron-chromium alloy produced by melting metallurgy.

DE 100 25 108 A1 discloses a high-temperature material comprising an iron alloy forming chromium oxide with up to 2% by weight of at least one oxygen-affine element from the group Y, Ce, Zr, Hf and Al, up to 2% by weight of an element M from the group Mn, Ni and Co, which forms a spinel phase of MCr₂O₄ type at high temperatures, up to 2% by weight of a further element from the group Ti, Hf, Sr, Ca and Zr, which increases the electrical conductivity of Cr-based oxides. The chromium content should lie within a concentration range between 12 and 28%. Areas of use for this high-temperature material are bipolar plates in a high-temperature fuel cell.

EP 1 298 228 A1 relates to a steel for a high-temperature fuel cell that has the following composition: not more than 0.2% C, not more than 1% Si, not more than 1% Mn, not more than 2% Ni, 15-30% Cr, not more than 1% Al, not more than 0.5% Y, not more than 0.2% REM and not more than 1% Zr, the remainder being iron and production-related impurities.

Low hot strength and inadequate creep strength at temperatures of 700° C. and above are common to these two alloys. Precisely in the range above 700° C. up to approximately 900° C., however, these alloys have excellent oxidation and corrosion resistance.

A creep-resistant ferritic steel, comprising precipitates of an intermetallic phase of Fe₂(M, Si) or Fe₇(M, Si)₆ type with at least one metallic alloying element M, which may be formed by the elements niobium, molybdenum, tungsten or tantalum, has become known from DE 10 2006 007 598 A1. The steel is preferably intended to be used for a bipolar plate in a fuel-cell stack.

EP 1 536 031 A1 discloses a metallic material for fuel cells, containing C≦0.2%, 0.02 to 1% Si, ≦2% Mn, 10 to 40% Cr, 0.03 to 5% Mo, 0.1 to 3% Nb, at least one of the elements from the group Sc, Y, La, Ce, Pr, Nd, Pm, Sn, Zr and Hf≦1, the remainder being iron and unavoidable impurities, wherein the composition is supposed to satisfy the following equation: 0.1≦Mo/Nb≦30.

EP 1 882 756 A1 describes a ferritic chromium steel, especially usable in fuel cells. The chromium steel has the following composition: C max. 0.1%, Si 0.1-1%, Mn max. 0.6%, Cr 15-25%, Ni max. 2%, Mo 0.5-2%, Nb 0.2-1.5%, Ti max. 0.5%, Zr max. 0.5%, REM max. 0.3%, Al max. 0.1%, N max. 0.07%, the remainder being Fe and melting-related impurities, wherein the content of Zr+Ti is at least 0.2%.

In comparison with DE 100 25 108 A1 and EP 1 298 228 A2, all of these alloys have better hot strength and elevated creep strength at temperatures of 700° C. and above, specifically due to formation of precipitates, which hinder the dislocation movements and thus plastic deformation of the material. In the case of DE 10 2006 007 598 A1, for example, these precipitates consist of a Laves phase, an intermetallic compound with the composition Fe₂(M, Si) or Fe₇(M, Si)₆, wherein M may be niobium, molybdenum, tungsten or tantalum. Therein a proportion by volume of 1 to 8%, preferably 2.5 to 5%, should be reached. However, there may also be other precipitates such as Fe-containing particles and/or Cr-containing particles and/or Si-containing particles, such as described, for example, in EP 1 536 031 A1, or there may be carbides containing Nb, W, Mo. It is common to all of these particles that they make deformation of the material difficult.

From the state of the art described above, it is known that small additions of Y, Zr, Ti, Hf, Ce, La and similar reactive elements can influence the oxidation resistance of Fe—Cr alloys very positively.

The alloys cited in DE 10 2006 007 598 A1, EP 1 536 031 A1 and EP 1 882 756 A1 are optimized for the application as interconnector plates for the high-temperature fuel cells: By use of a ferritic alloy containing 10 to 40% chromium, they have an expansion coefficient adapted as well as possible to the ceramic components anode and electrolyte.

Further requirements on the interconnector steel of a high-temperature fuel cell are, besides the creep strength already mentioned above, very good corrosion resistance, good conductivity of the oxide layer and little chromium volatilization.

The requirements on the reformers and the heat exchangers for the high-temperature fuel cell are the best possible creep strength, very good corrosion resistance and little chromium volatilization. The oxide for these components does not have to be conductive.

The requirements for components for the exhaust-gas line of a combustion engine, for example, or for steam boilers, superheaters, turbines and other parts of a power plant, are best possible creep strength and very good corrosion resistance. In these cases chromium volatilization does not cause any poisoning phenomena as in the fuel cell, and the protecting oxide does not have to be conductive for such components.

In DE 10 2006 007 598 A1, for example, the excellent corrosion resistance is achieved by formation of a chromium oxide top layer. By the fact that a spinel containing Mn, Ni, Co or Cu is additionally formed on the chromium oxide top layer, fewer volatile chromium oxides or chromium oxyhydroxides that poison the cathode are formed. By the fact that Si is stably bound in the Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phase, a nonconductive subsurface layer of silicon oxide is also not formed under the chromium oxide top layer. The corrosion resistance is further improved by the fact that the Al content is kept low and so the increase of the corrosion due to the internal oxidation of the aluminum is avoided. A small Ti addition additionally favors strengthening of the surface and thus prevents swelling of the oxide layer and the inclusion of metallic zones in the oxide layer, which increases the oxidation. In addition, the addition of oxygen-affine elements such as La, Ce, Y, Zr or the like further increases the corrosion resistance.

From the market, increased requirements are being imposed on products, necessitating elevated hot strength and creep strength together with an elongation of at least 18% at application temperature for avoidance of brittle failure together with at least equally good oxidation or corrosion resistance and a higher service temperature of the alloy, specifically while retaining acceptable deformability, measured as plastic deformation in the tension test with an elongation of >13% at room temperature.

Furthermore, the following investigation methods are used.

In a creep test, a specimen is subjected to a constant static tensile force at a constant temperature. For the purpose of comparability, this tensile force is expressed as an initial tensile stress relative to the initial cross section of the specimen. In the creep test, the time t_(B) until break—the time to break—of the specimen is measured in the simplest case. The test can then be performed without measurement of the elongation of the specimen in the course of the test. The elongation at break is then measured after the end of the test.

The specimen is mounted at room temperature in the creep-testing machine and heated to the desired temperature without loading by a tensile force. After reaching the test temperature, the specimen is maintained for one hour without loading for temperature equilibration. Thereafter the specimen is loaded with the tensile force and the test time begins.

The time to break can be taken as a measure of the creep strength. The longer the time to break is at a specified temperature and initial tensile stress, the greater the creep strength of the material is. The time to break and the creep strength decrease with increasing temperature and increasing initial tensile stress (see, for example, “Bürgel”, page 100).

The deformability is determined in a tension test according to DIN 50145 at room temperature. In the process, the offset yield strength R_(p0.2), the tensile strength R_(M) and the elongation at break are determined. The elongation A is determined on the broken specimen from the elongation of the original gauge length L₀:

A=(L _(u) −L ₀)/L ₀100%=ΔL/L ₀100%

Where L_(u)=gauge length after break.

Depending on gauge length, the elongation at break is denoted by subscripts:

A₅, gauge length L₀−5·d₀ or L₀=5.65·√S₀

A₁₀, gauge length L₀=10·d₀ or L₀=11.3··S₀

or, for example, A_(L=100), for the freely chosen gauge length L=100 mm. (d₀ initial diameter, S₀ initial cross section of the flat specimen)

The magnitude of the elongation A in the tension test at room temperature can be taken as a measure of the deformability.

The Laves phase(s) or the Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides can be made visible on a metallographic ground section by etching with V2A pickling fluid or electrolytic etching with oxalic acid. During etching with V2A pickling fluid, the grains or grain boundaries also are additionally etched visibly. Only particles with a size of approximately 0.5 μm and larger are visible by viewing in an optical microscope. Smaller particles may not be recognized, but are definitely present. Therefore metallography is used only in support of the explanation, while the efficacy of a measure is assessed more practically by the time to rupture or creep strength.

In the Manual of High-Temperature Materials Technology, Ralf Burgel, 3rd Revised Edition, Viehweg Verlag, December 2006, hereinafter referred to as “Bürgel”, the “Possible measures for increasing the creep strength of metallic materials” are presented on pages 196 to 199 and in Table 3.7.

The measures

-   -   “High melting point, face-centered cubic material”,     -   “High modulus of elasticity”,     -   “Material with low stacking fault energy”, cannot be used for         improvement of the cited parameters, since they necessitate a         change of material type, which is not possible here and also is         not the task.

The measures

-   -   “Solid solution hardening”     -   “Particle hardening”     -   “High particle volume fraction”     -   “Particles with small diffusion coefficients of the alloying         element in question”         have already been employed in DE 10 2006 007 598 A1 and/or in EP         1 536 031 A1 and/or in EP 1 882 756 A1.

The measures

-   -   “Particles with low solubility in the matrix”     -   “Coherent particles with low interfacial enthalpy relative to         the matrix”,         are not applicable for the precipitates under consideration.

Likewise, the measures

-   -   “Carbides or borides as grain-boundary precipitates; avoid         oxides and sulfides”,     -   “Add positively active grain-boundary elements in precisely         controlled dosage, for example, B, C, Zr, Ce”,     -   “Higher purity of the alloy”     -   “Add getter elements (for example, for S)”,     -   “High corrosion resistance”         have already been described in DE 10 2006 007 598 A1 and/or in         EP 1 536 031 A1 and/or in EP 1 882 756 A1.

The measures

-   -   “Small dendrite arm spacings in cast microstructures”,     -   “Small grain structure in the main loading direction”,     -   “Single crystal”     -   “Low density of components loaded by their own weight and         rotating” cannot be applied to this alloy type, or to the         production route, or the use.

For the task of improving the creep strength of the precipitation-hardened iron-chromium alloy, the measures

1) “Coarse grain microstructure”, 2) “Jaggedness of the grain boundaries due to precipitates”, 3) “Optimized heat treatment (adjust optimum particle diameter, eliminate segregations in cast microstructure, purposefully adjust possible grain boundary roughness)”, 4) “Avoid cold working”, are to be considered.

The task of the invention is to provide a method for production of a component made from a precipitation-hardened iron-chromium alloy, by means of which the high hot strength or creep strength of a precipitation-hardened ferritic alloy can be further increased compared with the state of the art while retaining acceptable deformability at room temperature.

It is also intended to provide a thermomechanically treated component/semifinished product consisting of an iron-chromium alloy, which can be used for achievement of high hot strength or creep strength while retaining acceptable deformability at room temperature.

Finally, it is intended that the component/semifinished product produced in this way can be used for specific technical applications in the temperature range above 550° C.

This task is accomplished on the one hand by a method for production of a component from an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, in that a semifinished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures≧the solution-annealing temperature, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in a second step mechanical working of the semifinished product in the range from 0.05 to 99% is performed and in a subsequent step Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form by the fact that the component made from the worked semifinished product is brought to an application temperature between 550° C. and 1000° C. by heating at 0.1° C./min to 1000° C./min.

This task is accomplished on the other hand by a method for production of a component from an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, in that a semifinished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures the solution-annealing temperature, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in a second step mechanical working of the semifinished product in the range from 0.05 to 99% is performed and in a subsequent step Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form by the fact that the worked semifinished product is subjected for a time between t_(aw), and t_(max) to a heat treatment in the temperature range between 550° C. and 1060° C. under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800° C. is quenched in the oven, wherein t_(min) and t_(max) are calculated according to the following formulas:

t _(min) =T _(a)·10^((6740/Ta−9.216)) and t _(max) =T _(a)·10^((17960/Ta−15.72)) where T _(a) =T+273.15,

and wherein the desired component is made before or after this heat treatment.

The times t_(min) and t_(max) are expressed in minutes and the heat treatment temperature T in ° C.

Advantageous further developments of the inventive method are to be inferred from the associated dependent claims pertaining to the method.

For the first step, the following temperature ranges and times are practical for solution annealing:

>1050° C. for longer than 6 minutes >1060° C. for longer than 1 minute

The resulting changes of the material characteristics are explained in more detail in the course of the further description.

Furthermore, the task is also accomplished by a metallic component or semifinished product consisting of the following chemical composition (in % by weight)

Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002-0.1% N 0.002-0.1% S max. 0.01%

Fe remainder as well as the usual melting-related impurities, which at the end of a thermomechanical treatment has a deformed microstructure, to the effect that Laves phase(s) is or are embedded in finely dispersed form in the microstructural dislocations of the microstructure, wherein, in a creep test with, for example, 35 MPa at 750° C. and an elongation of at least 18%, a time to break that exceeds the time to break of a coarse-grained, completely recrystallized microstructure by a factor of at least 1.5 is established in the microstructure.

A comparable result is achieved for creep tests with different stresses and temperatures, wherein the temperatures for the creep test preferably lie in the range between 500 and 1000° C.

Measures 1 to 4 described above will now be considered.

Surprisingly it has been found in this connection that, in contrast to measure 4 “cold deformation”, preworking followed by an adapted annealing treatment can bring about prolongations of the times to break of the specimen in the creep test that go more than 1.5 times, preferably more than 3 times beyond the times to break for a coarse-grained microstructure (measure 1).

Furthermore, it is proposed that, for the third step—the precipitation of the Laves phase(s)—the worked semifinished product or if applicable the component made therefrom, by a combination of heating at 0.1° C./min to 1000° C./min to a heat-treatment temperature between 550° C. and 1060° C. with subsequent heat treatment for a time between t_(min) and t_(max) at this temperature under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800° C. is quenched in the oven, after which the desired component is made if applicable, wherein t_(min) and t_(max) are calculated according to the following formulas:

t _(min) =T _(a)·10^((6740/Ta−9.216)) and t _(max) =T _(a)·10^((17960/Ta−15.72)) where T _(a) =T+273.15.

In addition, the possibility exists that, for the third step—the precipitation of the Laves phase(s)—the worked semifinished product or the component made therefrom is subjected for a time between t_(min) and t_(max) to a heat treatment in the temperature range between 550° C. and 1060° C. under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800° C. is quenched in the oven, after which the desired component is made if applicable, wherein t_(min) and t_(max) are calculated according to the following formulas: t_(min)=T_(a)·10^((6740/Ta−9.216)) and t_(max)=T_(a)·10^((17960/Ta−15.72)) where T_(a)=T+273.15 and then the component that has been made is brought by heating at 0.1° C./min to 1000° C./min to an application temperature between 550° C. and 1000° C.

According to a further concept of the invention, a semifinished product from an alloy of the following composition (in % by weight) is treated thermomechanically:

Cr 12 to 30% Mn 0.001 to 2.5% Nb 0.1 to 2% W 0.1 to 5% S 0.05 to 1% C 0.002 to 0.03% N 0.002 to 0.03% S max. 0.01%

Fe remainder as well as the usual melting-related impurities.

With the inventive method it is possible to produce semifinished products in the form of sheets, strips, bars, forgings, pipes or wire and to make components in the most diverse forms needed for the respective application.

It is of special advantage that only little or even no Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are still present in the semifinished product after solution annealing at temperatures≧the solution-annealing temperature, preferably ≧1050° C. for longer than 6 minutes or >1060° C. for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in the initial state before deformation.

The working of the semifinished product can take place by hot working. Alternatively, however, the forming can also be brought about by cold working.

In the first case, the semifinished product is hot-worked with a starting temperature ≧1070° C., wherein the last 0.05 to 95% of mechanical deformation is applied between 1000° and 500° C., advantageously the last 0.5 to 90% between 1000° C. and 500° C.

In the second case, the degree of cold working of the semifinished product is 0.05 to 99%, advantageously 0.05 to 95% or 0.05 to 90%.

As a further concept according to the invention, it is proposed that the mechanical working of the semifinished product be 20 to 99% and then the worked semifinished product be subjected for a time between t_(min) and t_(max) to a heat treatment in the temperature range between 950° C. and 1060° C. under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, after which the desired component be made, wherein t_(min) and t_(max) are calculated according to the following formulas:

t _(min) T _(a)·10^((6740/Ta−9.216)) and

t _(max) =T _(a)·10^((17960/Ta−15.72)) where T _(a) =T+273.15

with t_(min) and t_(max) in minutes and the heat-treatment temperature T in degrees Celsius. If the already indicated alloy is used as an interconnector for a solid oxide fuel cell, then a content of 0.001-0.5% aluminum is advantageous.

For other areas of use, such as, for example, in the reformer or heat exchanger for the fuel cell, for which no conductive oxide layer is necessary, a content of 2 to 6% aluminum is advantageous, since then a closed aluminum oxide layer can form, which once again has a much slower growth rate compared with a chromium oxide layer and additionally has much less chromium oxide volatilization than a chromium-manganese spinel.

For the areas of use that neither need a conducive oxide nor have special requirements on chromium volatilization, both variants may be considered. In this connection, it is to be kept in mind in particular that the processability and weldability of the alloy deteriorate with increasing aluminum content and so higher costs are incurred. Therefore, when an oxide layer consisting of a chromium oxide and a chromium-manganese spinel, adequate oxidation resistance can be assured by use of 0.001-0.5% aluminum. If greater oxidation resistance is necessary, as is assured, for example, by the formation of an aluminum oxide layer, a content of 2.0-6.0% aluminum is advantageous. These two alloy variants can be used, for example, as components for the exhaust-gas line of a combustion engine or for steam boilers, superheaters, turbines and other parts of a power plant.

A preferred aluminum range is in particular the range from 2.5% to 5.0%, which is still characterized by good processability.

In the already indicated alloy, the following elements may be additionally used individually or in combination:

La 0.02 to 0.3% Ti 0.01 to 0.5% Mg 0.0001 to 0.07% Ca 0.0001 to 0.07% P 0.002 to 0.03% Ni/Co/Cu 0.01 to 3% B up to 0.005%.

The contents of the elements that can be additionally introduced in the alloy may be adjusted as follows: Mg 0.0001 to 0.05%, Ca 0.0001 to 0.03%, P 0.002 to 0.03%.

Furthermore, the alloy (in % by weight) may contain one or more of the elements Ce, La, Pr, Ne, Sc, Y, Zr or Hf in contents of 0.02-0.3%.

If necessary, the alloy (in % by weight) may contain one or more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf in contents of 0.02-0.2%.

To achieve the desired effects, the Nb content is 0.3 to 1.0% and the Si content 0.15 to 0.5%.

If necessary, the element tungsten may be replaced entirely or partly by at least one of the elements Mo or Ta.

If necessary, the alloy may also even contain max. 0.2% V and/or max. 0.005% S. In this case the oxygen content should not be greater than 0.01%.

If necessary, the alloy may also even contain max. 0.003% boron.

Furthermore, the alloy should have a maximum of 0.01% of the following elements respectively: Zn, Sn, Pb, Se, Te, Bi, Sb.

Components/semifinished products that on the one hand consist of the cited alloy composition and on the other hand have been produced by the inventive method may preferably be used as interconnector in a fuel cell or as material in a component, such as a reformer or a heat exchanger in an ancillary aggregate of the fuel cell.

Alternatively, the possibility also exists of using the component/semifinished product produced according to the inventive method or the alloy itself as a structural element in the exhaust-gas line of a combustion engine or for steam boilers, superheaters, turbines and other parts of a power plant or in the chemical process industry.

By means of the inventive method, Laves phases, by virtue of the thermomechanical treatment, can be precipitated purposefully and in fine dispersion at the dislocations of the microstructure in alloys produced by melting metallurgy.

The details and the advantages of the invention will be explained in more detail in the following examples.

In the following, the inventive method steps will be subjected to closer examination.

The first step for the thermomechanical treatment of an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides must be annealing above the solution annealing temperature, so that the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are dissolved and are available for precipitation in the subsequent thermomechanical treatment. The solution annealing temperature is alloy-specific, but preferably lies above 1050° C. for a period of longer than 6 minutes or above 1060° C. for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water. The exact temperature control above this solution annealing temperature is not determining for the characteristics. The annealing may be carried out in air or under protective gas. It should lie under the melting temperature, preferably <1350° C. Also, for cost reasons, the annealing times should preferably be <24 hours, but may also be longer depending on performance. The solution annealing follows quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, during which only little Laves phase is newly formed.

In addition, care is to be taken that, especially for thicker-walled components, all parts of the component reach the required minimum annealing time at the specified temperature. This is to be considered in the determination of the starting point of the annealing time.

In a second step, an elevated dislocation density must be introduced into the material. Elevated dislocation densities have worked microstructure or recovered microstructure, wherein the dislocations there are arranged at small-angle grain boundaries.

The second step must therefore be working, so that the dislocations are introduced into the material, which then, in the subsequent annealing treatment, ensure a homogeneous dispersion of the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides.

This deformation may be cold working, but also hot working, wherein it must be ensured during hot working that the microstructure is not already recrystallized during rolling. This is achieved by restricting the deformation range for the last working and the temperature at which this is carried out. For deformations above 1000° C., the material already tends to recrystallization or recovery during working, so that the working must preferably be carried out below 1000° C. At temperatures below 500° C., exist in the range of the embrittlement that occurs in ferrites at 475° C. There this has a smaller elongation and an elevated working resistance, which makes working less advantageous and reduces the economic benefit.

Precipitates smaller than a certain size are less effective (see, for example, “Bürgel”, page 141). Therefore the dislocation density generated by the deformation should not be too high, since then very many precipitates are indeed formed but are too fine, and the excess dislocations can move freely and in this way the preworking becomes harmful. This means that preferably the greatest deformation is 90% for the part of the hot working ≦1000° C. and 90% for the cold working.

During working in the range of 20 to 99%, annealing between 950° C. and 1050° C. may cause recovery of the microstructure. Thereby the dislocation density is reduced, so that the positive effect on the dispersion of the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides is established again.

The one possibility of introducing the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides into the worked material is to make the needed components from the semifinished product and then to bring the component that has been made to the application temperature between 550° C. and 1000° C. by heating at 0.1° C./min to 1000° C./min. During heating, the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are precipitated as a fine dispersion in the microstructure. The fine dispersion is generated by nucleation in the lower temperature range, followed by some growth of the nuclei at the higher temperatures. Therefore the heating rate should not be slower than 1000° C./min, because otherwise the time for this process is too short. Heating rates slower than 0.1° C./min are uneconomical.

A second possibility is a separate heat treatment of the material. For this purpose, the worked semifinished product/component is subjected for a time between t_(min) and t_(max) to a heat treatment in the temperature range between 550° C. and 1060° C. under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800° C. is quenched in the oven, wherein

t _(min) =T _(a)·10^((6740/Ta−9.216)) and

t _(max) =T _(a)·10^((17960/Ta−15.72)) where T _(a) =T+273.15,

with indication of t_(min) and t_(max) in minutes and heat treatment temperature T in ° C. In this connection, the desired component can be made before or after this heat treatment.

In the annealing steps, care is to be taken that, especially for thicker-walled semifinished products/components, all parts of the component reach the required minimum annealing time at the specified temperature. This is to be considered in determination of the starting point of the annealing time. Likewise, care is to be taken that no region of the semifinished product/component exceeds the required maximum annealing time.

Times shorter than t_(min) are not sufficient for formation of the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides. For times longer than t_(max), the danger exists of too great coarseness of the precipitates, whereby the particles can no longer contribute markedly to the creep strength. For times longer than t_(max), the possibility exists in the upper temperature range of 550° C. and 1060° C. that a recovered microstructure will be formed, which certainly can still be effective. However, with increasing recovery, the dislocation density is further reduced, so that the dispersion of the precipitates becomes increasingly inhomogeneous and the positive effect on the creep strength ultimately vanishes. In addition, at the lower temperatures in the range of 550° C. and 1060° C., times longer than t_(max) are uneconomical.

The annealing step may be carried out under protective gas (argon, hydrogen and similar atmospheres with reduced oxygen partial pressure). For economic reasons, the quenching step is carried out in stationary protective gas or air, moving (blown) protective gas or air or in water, while oven quenching should be avoided in particular for temperatures above 800° C. but is also possible at temperatures <800° C.

The oxidation resistance and the thermal expansion coefficient of the material are determined via the chromium content. The oxidation resistance of the material is based on the formation of a closed chromium oxide layer. Below 12%, iron-containing oxides, which impair the oxidation resistance, are formed increasingly, especially at higher operating temperatures. The chromium content is therefore adjusted to >12%. Above 30% chromium, the processability of the material and its usability are impaired by increased formation of embrittling phases, especially the sigma phase. The chromium content is therefore limited to ≦30%. The expansion coefficient decreases with increasing chromium content.

Especially for use in a fuel cell, therefore, the expansion coefficient can be adjusted in a range that matches the ceramics in the fuel cell. These are chromium contents around 22 to 23%. For other applications, however, for example for reformers or in power plants, this restriction does not exist.

The addition of manganese brings about formation of a chromium-manganese spinel on the chromium oxide layer, which is formed on the material at low aluminum contents below 2%. This chromium-manganese spinel reduces the chromium volatilization and improves the contact resistance. A manganese content of at least 0.001% is necessary for this. More than 2.5% manganese impairs the oxidation resistance by formation of a very thick chromium-manganese spinel layer.

Niobium, molybdenum, tungsten or tantalum can participate in the formation of precipitates in iron-containing alloys, such as, for example, carbides and/or the M in the Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases. Molybdenum, tungsten or tantalum are also good solid-solution hardeners and thus contribute to the improvement of the creep strength. In this connection the lower limit is determined in each case by the fact that a certain content must be present in order to be effective, while the upper limit is determined by the processability. Thus the preferred ranges are, for

Nb 0.1-2% W: 0.1-5%

W may also be replaced entirely or partly by Mo and/or Ta: 0.1-5%.

Silicon can participate in the formation of precipitates in iron-containing alloys, for example in the Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases. It favors the increased precipitation and stability of these Laves phases and in this way contributes to the creep strength. During formation of the Laves phases, it is completely bound in these. Thus the formation of a silicon dioxide layer no longer takes place under the chromium oxide layer. At the same time, the incorporation of M in the oxide layer is reduced, whereby the negative influence of M on the oxidation resistance is prevented. At least 0.05% Si must be present for the desired effect to occur. If the content of Si is too high, the negative effect of the Si may reappear. The Si content is therefore limited to 1%.

Aluminum in contents below 1% impairs the oxidation resistance, since it leads to internal oxidation. However, an aluminum content higher than 1% leads to formation, under the chromium oxide layer, of an aluminum oxide layer, which is not electrically conductive and thus reduces the contact resistance. Therefore the aluminum content is limited to ≦0.5% when a chromium oxide former is desired or its oxidation resistance is sufficient. An example of this is, for example, for the application as interconnector plate. However, a certain aluminum content of at least 0.001% is necessary for deoxidation of the melt. If no conductive oxide is necessary and at the same time the requirement of much higher oxidation resistance than is given by a chromium oxide layer is still required, the alloy may form a closed aluminum oxide layer by a content of aluminum of at least 2% (DE 101 20 561). Aluminum contents above 6.0% lead to processing problems and thus to increased costs.

Carbon leads to carbide precipitates and thus contributes to the creep strength. The carbon content should be <0.1%, in order not to impair the processability. However, it should be >0.002%, so that an effect can occur.

The nitrogen content should be 0.1% maximum, in order to avoid formation of nitrides, which impair the processability. It should be higher than 0.002%, in order to assure the processability of the material.

The contents of sulfur should be made as low as possible, since this interface-active element impairs the oxidation resistance. A maximum of 0.01% S is therefore stipulated.

Oxygen-affine elements such as Ce, La, Pr, Ne, Sc, Y, Zr, Hf improve the oxidation resistance by reducing the oxide growth and improving the adherence of the oxide layer. A minimum content of 0.02% of one or more of the elements Ce, La, Pr, Ne, Sc, Y, Zr, Hf is practical in order to obtain the oxidation-resistance-increasing effect of the Y. For cost reasons, the upper limit is set at 0.3% by weight.

As with every oxygen-affine element, titanium is bound in the oxide layer during oxidation. In addition, it also causes internal oxidation. However, the resulting oxides are so small and finely dispersed that they cause hardening of the surface and thus prevent swelling of the oxide layer and inclusion of metallic zones during the oxidation (see DE 10 2006 007 598 A1). These swellings are unfavorable, since the resulting cracks cause an increase of the oxidation rate. Thus Ti contributes to the improvement of the oxidation resistance. For effectiveness of the Ti content, at least 0.01% Ti must be present, but not more than 0.5%, since this does not improve the effect further but increases the costs.

The content of phosphorus should be lower than 0.030%, since this interface-active element impairs the oxidation resistance. Too low P content increases the costs. The P content is therefore ≧0.002%.

The contents of magnesium and calcium are adjusted in the spread ranges of 0.0001 to 0.05% by weight and 0.0001 to 0.03% by weight respectively.

It has been found that cobalt contents of 3% and higher impair the oxidation resistance. For cost reasons, the lower limit is set at 0.01% by weight. For nickel and copper, the same applies as for cobalt.

Boron is limited to max. 0.005%, since this element reduces the oxidation resistance.

The subject matter of the invention will now be explained in more details on the basis of exemplary embodiments.

The analyses of the batches used for the following examples are presented in Table 1. These batches were melted in the arc furnace in an amount of approximately 30 metric tons, thereafter cast in a ladle and subjected to a decarburization and deoxidation treatment as well as to a vacuum treatment in a VOD system and cast to ingots. These were then hot-rolled and, depending on final thickness, cold-rolled with intermediate annealing steps. After the hot-rolling, the oxide layer was removed by pickling.

A material with an analysis as indicated in Table 1 precipitates mainly Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases and, in much smaller contents, carbides.

EXAMPLE 1

In this example, material from the batch 161061 listed in Table 1 was hot-rolled to 12 mm thick sheet after solution annealing above 1070° C. for a period of longer than 7 minutes followed by quenching in stationary air, wherein the mechanical working was begun with a start temperature >1070° C. and the last 78% of mechanical deformation was applied by rolling between 500° C. and 1000° C.

FIG. 1 shows the typical appearance of a microstructure deformed in this way. In the microsections etched by means of electrolytic etching with oxalic acid, it can be clearly seen that only little Laves phase has been precipitated in microscopically visible form.

When the material formed in this way is then annealed at 1075° C. for 20 minutes with quenching in stationary air, a microstructure is obtained with only few precipitates of Laves phase and a grain size of approximately 137 μm (FIG. 2), which is a typical coarse-grained microstructure.

When a creep test as described above is performed on this material with an initial stress of 35 MPa at a temperature of 750° C., the specimen breaks after 12.8 hours at an elongation A of 69.8%. (Table 2). At room temperature, the material annealed in this way has an elongation of 35%, which is a very good value for a ferrite.

If, in contrast, from the hot-rolled material, which is synonymous with preworking, a specimen for a creep test is made as simulation for a component and this is then heated at approximately 60° C./minute to an application temperature of 750° C. and then a creep test is performed with an initial stress of 35 MPa at a temperature of 750° C., the specimen surprisingly breaks only after 255 hours at an elongation A of 29%, which means a prolongation of the time to break by a factor of 20. Making the component is very easily possible, since the hot-worked condition, as was described above, has an elongation of 19% in the tension test at room temperature, which is a good value and makes the material readily processable.

This example clearly shows that the microstructure with the preworking and the coarse-grained microstructure is superior with respect to time to break or creep strength, which contradicts the state of the art as described in “Bürgel” pages 196 to 199 Table 3.7.

EXAMPLE 2

In this example, annealing steps were performed on the hot-rolled material from Example 1 for 20 minutes each between 600° C. and 1000° C. or for some temperatures also for 240 or 1440 minutes (see Table 3 for t_(min) and t_(max) according to Equation 1 and 2) in air, followed by quenching in stationary air. After the heat treatment, specimens were made from the sheet and then the creep test was performed with a stress of 35 MPa at 750° C. as described above. The results are compiled in Table 3.

After 20 minutes at 1000° C., a time to break of only 10.4 hours is reached at an elongation A of 79.5%. After 20 minutes at 600° C. to 950° C., a time to break of longer than 100 hours, at least 7 times longer, is achieved at an elongation A of greater than 22.7%. The longest time to break for annealing steps of 20 minutes is achieved at 850° C., with 564 hours. The longest time to break for annealing steps of 240 minutes is achieved at 800° C., with 396 hours. After 1440 minutes at 700° C., a time to break of 645 hours is even reached. FIG. 3 shows the microstructure after the various annealing steps for 20 minutes. The microstructures in FIG. 3 are not globularly recrystallized. Up to 850° C. (the maximum of the time to break), the microstructure has the typical appearance of a deformed microstructure. Starting from approximately 900° C., recovery can be clearly recognized, but this means that the dislocation density is still increased compared with a globularly recrystallized microstructure. In a recovered microstructure, the dislocations have become partly reordered at small-angle grain boundaries. This has an effect similar to that of preworking. In the microsections etched by means of electrolytic etching with oxalic acid, it can be clearly recognized that the Laves phase is precipitated in microscopically visible form starting from approximately 750° C., wherein it is precipitated increasingly more densely and more homogeneously up to 850° C. (the maximum of the time to break). From approximately 900° C. on, small-angle grain boundaries or grain boundaries can also be recognized markedly besides the precipitates in the grain, thus assuring jaggedness of the small-angle grain boundaries or grain boundaries, which corresponds to measure 2 for increasing the creep strength (see above). At 1000° C., very large grains have recognizably formed due to advancing recovery, such that the dislocation density is greatly reduced and so no further increase of the time to break occurs. The maximum of the time to break occurs in the deformed microstructure with dense homogeneously precipitated Laves phases.

At room temperature, the sheet annealed for 20 minutes at all temperatures between 600° C. and 950° C. has an elongation of at least 13%, which is still to be regarded as satisfactory for a ferritic alloy and makes the material processable. The elongation is smallest in the range of 700° C. to 800° C. and is improved at the lower or higher annealing temperatures respectively, because at the lower temperatures Laves phase is certainly already precipitated but is not yet microscopically visible and therefore has a smaller proportion by volume, but in return is very finely dispersed. At the higher temperatures, a larger proportion by volume is precipitated, but in return is somewhat coarser and recognizable at the small-angle grain boundaries and grain boundaries.

The annealing at 1000° C. for an annealing time of 20 minutes exceeds t_(max)=19.6 minutes. Thus it is not in the range of the invention and is used as reference. Also, the time to break is only 10.4 hours. The annealing time of 20 minutes at the temperatures between 600° C. and 950° C. lies in the inventive range between t_(min) and t_(max). Accordingly, the time to break was clearly prolonged according to the invention by more than a factor of 7 compared with the coarse-grained, globularly recrystallized condition from Example 1, which is obtained after annealing at 1075° C./20 minutes followed by quenching in stationary air.

EXAMPLE 3

In this example, material from the batch 161061 listed in Table 1 was hot-rolled to 12 mm thick sheet after solution annealing above 1070° C. for a period of longer than 7 minutes followed by quenching in stationary air, wherein the working was begun with a start temperature >1070° C. and the last 60% of mechanical working was applied by rolling between 1000° C. and 500° C.

If the sheet worked in this way is then annealed industrially in the continuous furnace at 920° C. for 28 minutes in air and quenched in stationary air, a tension specimen made from this material has a time to break of 391 hours at an elongation A of 38% (Table 4) in the creep test with an initial stress of 35 MPa at a temperature of 750° C. The microstructure is not globularly recrystallized but instead is recovered. It has precipitates in the grain and at the small-angle grain boundaries or grain boundaries (FIG. 4). The time to break is 30 times the time achieved in Example 1 after annealing at 1075° C. for 20 minutes with a globularly recrystallized coarse-grained microstructure with a grain size of 137 μm. The annealing at 920° C. for an annealing time of 28 minutes lies in the inventive range between t_(min)=0.32 minutes and t_(max)=162.6 minutes.

At room temperature, the sheet treated in this way has a very good elongation of 18%, an offset yield strength of 475 MPa and a tensile strength of 655 MPa (see Table 4), which makes the material readily workable.

EXAMPLE 4

In this example, material from batch 161061 and batch 161995 was cold-rolled to 1.5 mm thick sheet after solution annealing at above 1070° C. for a period longer than 7 minutes followed by quenching in blown air and hot rolling as well as removal of the oxide layer, wherein cold working of 53% was applied. Subsequently annealing at 1050° C. was carried out for 3.4 minutes under protective gas in the continuous furnace with subsequent quenching in the cold stream of protective gas. Thereafter both batch 161061 (FIG. 5) and batch 161995 exhibit a recovered microstructure with elongated grains (FIG. 7) and precipitation of Laves phase, although much less than recognizable in FIG. 4. Thereafter part of the material was annealed once again at 1050° C. for 20 minutes under air with subsequent quenching in stationary air. After this, both batches were globularly recrystallized, batch 161061 with a grain size of 134 μm (FIG. 6) and batch 161995 with a grain size of 139 μm. Only slight precipitated Laves phase can still be found.

Tables 5a and 5b show the results of the creep tests and of the tension tests at room temperature. After the annealing at 1050° C. for 3.4 minutes, batch 161061 has a time to break of 25.9 hours at an elongation A of 50% in a creep test at 750° C. with an initial load of 35 MPa, and after additional annealing at 1050° C. for 20 minutes, which produces very coarse grain, a time to break of only one third, 7.9 hours, at an elongation A of 83%.

Similarly, batch 161995 has a time to break of 33.5 hours 89% in a creep test at 750° C. with an initial load of 35 MPa, and after additional annealing at 1075° C. for 20 minutes, which produces very coarse grain, a time to break of only one third, 7.9 hours, at an elongation A of 92%. The elongation of 28% in the tension test at room temperature for batch 161061 for 1050° C. and 3.4 minutes of annealing time and of 26% for batch 161995 is very good for a ferrite, which makes the material very readily workable. For the coarse-grained structure, it is even higher, with 31% for batch 161061 and 29% for batch 161995.

This shows the influence of the annealing time at temperatures around 1050° C., In short-time annealing steps of a few minutes, dislocations (deformation) and adequate Laves phase are present in the material, which in this example has the consequence of a 3 to 4 times longer time to break in the creep test. For longer annealing steps, the Laves phase is sufficiently dissolved, as batch 161061 shows, and the microstructure recrystallizes globularly with correspondingly short times to break in the creep test.

The annealing at 1050° C. for 20 minutes lies with an annealing time of 20 minutes above t_(max)=6.0 minutes. Thus it does not fall within the range of the invention and is used as reference, just as the annealing at 1075° C. for 20 minutes. The annealing at 1050° C. for 3.4 minutes lies with an annealing time of 3.4 minutes in the inventive range between t_(min)=0.32 minutes and t_(max)=6.0 minutes and according to the invention exhibits a clearly improved time to break in the creep test.

EXAMPLE 5

In this example, material from the batch 161061 was hot-rolled to 12 mm thick sheet after solution annealing above 1070° C. for a period of longer than 7 minutes followed by quenching in stationary air, wherein the working was begun with a start temperature >1070° C. and the last 70% of mechanical deformation was applied by rolling between 1000° C. and 500° C.

When the material worked in this way is then subjected to solution annealing at 1075° C. for 22 minutes with quenching in stationary air, a very coarse-grained microstructure is obtained with only few precipitates of Laves phase and a grain size of approximately 134 to 162 μm (FIG. 9). When a creep test is performed on this material with an initial stress of 40 MPa at a temperature of 700° C., the specimen breaks after 228 hours at an elongation A of 51%. (Table 6) When the creep test is performed at 60 MPa, the specimen breaks after 8.1 hours, at an elongation A of 43%. At room temperature, the material annealed in this way has an elongation of 35%, which is a very good value for a ferrite.

If the material solution annealed at 1075° C. for 22 minutes is additionally subjected to annealing for 4 hours at 700° C. with subsequent quenching in stationary air, Laves phase dispersed in the microstructure is precipitated. (FIG. 10). When the creep test is then performed at 700° C. with an initial stress of 40 MPa, the specimen already breaks after 104 hours, at an elongation A of 72.6%, therefore a much shorter time than after the solution annealing at 1075° C. for 22 minutes. When the creep test is performed at 60 MPa, the specimen breaks after 6.3 hours, at an elongation A of 63%, therefore also after substantially shorter time than after the solution annealing at 1075° C. for 22 minutes.

This is the proof that the precipitation of the Laves phase(s) must take place in a microstructure with elevated dislocation density, therefore in a worked or recovered microstructure, in order to achieve prolongation of the time to break. Precipitation in a solution annealed microstructure has exactly the opposite effect, namely shortening of the time to break. The cause of this is the more homogeneous dispersion of very fine precipitates in the case of precipitation in a microstructure with elevated dislocation density, or in other words a deformed or recovered microstructure, in comparison with precipitation in a dislocation-poor coarse-grained microstructure.

EXAMPLE 6

In this example, as in Example 2, annealing steps were performed on the hot-rolled material from Example 1 for 20 minutes each between 750° C. and 1000° C. or for some temperatures also for 120 minutes, 240 minutes, 480 minutes, 960 minutes, 1440 minutes or 5760 minutes (see Table 7 for t_(min) and t_(maX) according to Equation 1 and 2) in air, followed by quenching in stationary air. After the heat treatment, specimens were made from the sheets and then the creep test was performed with a stress of 40 MPa at 750° C. as described above. The higher stress in comparison with Example 2 was chosen for shortening of the test time. The objective was to find heat-treatment times suitable for the annealing steps. The results are compiled in Table 7.

After 20 minutes at 1000° C., a time to break of only 8.8 hours is reached at an elongation A of 78.7%. In Example 2, after 20 minutes at 1000° C. and a creep test at 750° C. and 35 MPa, a time to break comparable with that after solution annealing at 1075° C. for 20 minutes with quenching in stationary air was reached, and so this value can be taken as reference for the time to break of the solution annealed condition. Inventive variants should also exceed this break time once again by a factor of at least 1.5.

After 20 minutes at 750° C. to 900° C., a time to break of longer than 100 hours, at least 10 times longer, is achieved at an elongation A of greater than 27%. The longest time to break for annealing steps of 20 minutes is achieved at 850° C. with 296 hours. The longest time to break for annealing steps of 120 minutes is achieved at 800° C. with 227 hours. The longest time to break for annealing steps of 240 minutes is achieved at 750° C. with 182 hours, but in this connection no value exists for 700° C. The longest time to break for annealing steps of 480 minutes is achieved at 800° C. with 169 hours. For 960 minutes, only one time to break was determined, for 750° C., with a value of 139 hours at an elongation of 24.2%. After 1440 minutes and 5760 minutes at 750° C. and 800° C., only times to break clearly shorter than the maximum times to break achieved at these temperatures are still achieved. At 800° C., for example, the time to break after a treatment time of 480 minutes drops from 169 hours to a value of 46 hours after a treatment time of 1440 minutes, to a value of 17.5 hours after a treatment time of 5760 minutes, although this is still in the inventive range. A further prolongation of the treatment time should shorten the time to break further, so that t_(max) of 7059 minutes is logically somewhat longer than the time of 5750 minutes. All elongations for the heat treatment temperatures from 750 to 900° C. and times from 20 minutes to 5760 minutes lie between 24.2% and 43% and therefore are larger than 18%, as required, in order to avoid brittle failure. For the microstructure after 20 minutes of annealing, what was said in Example 2 is applicable, since the annealing steps are the same. Even at the higher stress of 40 MPa in the creep test, the maximum of the time to break occurs in the deformed microstructure, with dense homogeneously precipitated Laves phase.

At room temperature, the sheet annealed for 20 minutes at all temperatures between 600° C. and 900° C. in Example 2 has an elongation of at least 13%, which is still to be regarded as satisfactory for a ferritic alloy and makes the material processable.

EXAMPLE 7

In this example the 1.5 mm thick material of batch 161995, which was annealed after cold working of 53% at 1050° C. for 3.4 minutes under protective gas in the continuous furnace with subsequent quenching in the stream of cold protective gas, was used once again. In the same manner, 2.5 mm thick material from batch 161995 was produced by annealing it, after cold working of 40%, at 1050° C. for 2.8 minutes under protective gas in the continuous furnace with subsequent quenching in the stream of cold protective gas. Even the 2.5 mm thick material then exhibits a recovered microstructure with elongated grains (FIG. 11), just as the material from Example 4 in FIG. 7, and precipitation of Laves phase, albeit clearly less than recognizable in FIG. 4. Part of the material was then annealed once more at 1050° C. for 10 minutes under air with subsequent quenching in stationary air. After this the material had globularly recrystallized structure, with a grain size of 108 μm. Only little precipitated Laves phase is still to be found. The material with 1050° C./2.8 minutes and the material with 1075° C./10 minutes as the last heat treatment was then rolled with degrees of working between 2.8 and 40%. Thereafter creep tests were performed at 750° C. and 35 MPa and tension tests were performed at room temperature. The results are summarized in Table 8.

After the annealing at 1050° C. for 3.4 minutes, batch 161995, in a creep test at 750° C. with an initial load of 35 MPa, had a time to break of 33.5 hours at an elongation A of 89% and, after the additional annealing at 1050° C. for 10 minutes, which produces very coarse grain, it had a time to break amounting to only one third, 10.8 hours, at an elongation A of 50.4%.

If, from the material worked after 1050° C./2.8 minutes, a specimen for a creep test is made as simulation for a component and this is then heated at approximately 60° C./minute to an application temperature of 750° C. and then a creep test is performed with an initial stress of 35 MPa at a temperature of 750° C., the elongation at break for degrees of working between 5 and 40% drops to values around the 10 hours with elongations at break greater than 45%.

If, in contrast, from the material formed after 1050° C./10 minutes, a specimen for a creep test is made as simulation for a component and this is then heated at approximately 60° C./minute to an application temperature of 750° C. and then a creep test is performed with an initial stress of 35 MPa at a temperature of 750° C., the elongation at break for degrees of working between 2.9 and 40% increases to values between 49 and 137 hours, which means an increase by more than a factor of 4 in the time to break compared with the material worked after 1050° C./2.8 minutes, wherein a maximum occurs at 10% and the elongations at break lie between 18.9 and 60%.

From 10% degree of working on, however, the elongation at break in the tension test at room temperature becomes smaller than 8%, so that the material is increasingly more poorly processable. In other words, preferred degrees of forming for cold shaping lie between 0.05 and 10%. Thus this example shows that the increase of the elongation at break after working does not occur if the annealing before working was carried out at too low temperatures or for too short times. (Here at 1050° C. for 2.8 minutes) The increase of the elongation at break occurred after annealing was carried out at >1050° C. for >6 minutes (here at 1075° C. for 7 minutes).

The titles/descriptions of the tables/figures are reproduced as follows:

-   Table 1 Composition of the investigated alloy (all values in % by     weight) -   Table 2 Results of the creep tests at 750° C. with 35 MPa and of the     tension tests at room temperature for the hot rolling and the heat     treatments in Example 1 for a 12 mm thick sheet. (R: reference     according to the state of the art, I: according to the invention) -   Table 3 Results of the creep tests at 750° C. with 35 MPa and of the     tension tests at room temperature for the hot rolling from Example 1     and the heat treatment from Example 2 for a 12 mm thick sheet. (R:     reference according to the state of the art, I: according to the     invention) -   Table 4 Results of the creep tests at 750° C. with 35 MPa and of the     tension tests at room temperature for Example 3 for a 12 mm thick     sheet. (R: reference according to the state of the art, I: according     to the invention) -   Table 5 Results of the creep tests at 750° C. with 35 MPa and of the     tension tests at room temperature for Example 4 for a 1.5 mm thick     strip. (R: reference according to the state of the art, I: according     to the invention) -   Table 6 Results of the creep tests at 700° C. and of the tension     tests at room temperature for Example 5 on 12 mm thick sheet (R:     reference according to the state of the art, I: according to the     invention) -   Table 7 Results of the creep tests at 750° C. with 40 MPa and of the     tension tests at room temperature for the hot rolling and the heat     treatments in Example 6 for a 12 mm thick sheet. (R: reference     according to the state of the art, I: according to the invention) -   Table 8 Results of the creep tests at 750° C. with 35 MPa and of the     tension tests at room temperature for Example 7 for 1.5 mm to 2.5 mm     thick strip from batch 161995. (R: reference according to the state     of the art, I: according to the invention)

FIG. 1 Microstructure of the hot-worked material in Example 1

FIG. 2 Microstructure of the hot-worked material in Example 1 after annealing at 1075° C. for 20 minutes and quenching in stationary air, grain size 137 μm.

FIG. 3 Microstructure of the material in Example 2 after annealing between 600° C. and 1000° C. for 20 minutes in each case and quenching in stationary air.

FIG. 4 Microstructure of the material in Example 3 after annealing at 920° C. in the continuous furnace in air with subsequent quenching in stationary air for 20 minutes in each case and quenching in stationary air. (etching with V2A pickling fluid)

FIG. 5 Microstructure of batch 161061 in Example 4 after annealing at 1050° C./3.4 minutes under, protective gas in the continuous furnace with quenching in the stream of cold protective gas.

FIG. 6 Microstructure of batch 161061 in Example 4 after annealing at 1050° C./3.4 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas and annealing at 1050° C./20 minutes under air with subsequent quenching in stationary air, grain size 134 μm (etching with V2A pickling fluid)

FIG. 7 Microstructure of batch 161995 in Example 5 after annealing at 1050° C./3.4 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas.

FIG. 8 Microstructure of batch 161995 in Example 5 after annealing at 1050° C./3.4 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas and annealing at 1075° C./20 minutes under air with subsequent quenching in stationary air, grain size 139

FIG. 9 Microstructure of the hot-worked material in Example 5 after annealing at 1075° C. for 22 minutes and quenching in stationary air, grain size 134 μm to 162 μm

FIG. 10 Microstructure of the hot-worked material in Example 5 after annealing at 1075° C. for 22 minutes followed by quenching in stationary air and subsequent annealing at 700° C. for 4 hours followed by quenching in stationary air. Grain size 136 μm.

FIG. 11 Microstructure of batch 161995 in Example 7 after annealing at 1050° C. for 2.8 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas.

FIG. 12 Microstructure of batch 161995 in Example 7 after annealing at 1050° C. for 2.8 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas, followed by annealing at 1075° C. for 10 minutes under air with subsequent quenching in stationary air. Grain size 108 μm.

TABLE 1 Batch Batch Element 161061 161995 C 0.007 0.009 S <0.002 <0.002 N 0.015 0.018 Cr 22.9 22.6 Ni 0.30 0.22 Mn 0.43 0.43 Si 0.21 0.24 Mo 0.02 0.02 Ti 0.07 0.06 Nb 0.51 0.49 Cu 0.02 0.02 Fe Remainder Remainder P 0.014 0.017 Al 0.02% 0.019 Mg 0.0006 <0.01 Pb <0.001 <0.001 Sn <0.01 <0.01 Ca 0.0002 0.01 V 0.05 0.02 Zr <0.01 <0.01 W 1.94 1.97 Co 0.04 0.02 La 0.08 0.05 Ce <0.01 O 0.004

TABLE 2 Creep test with 35 Tension test at room temperature MPa at 750° C. Offset yield Tensile Time t_(B) _(—) to Elongation strength strength Elongation Heat treatment break in hours A in % Rp0.2 RM in MPa A 5 in % R 1075° C./20 minutes 12.8 69.8 359 494 35 Hot-worked 511 604 19 I Component + heating at 255 29.0 60° C./minute to application temperature 750° C.

TABLE 3 Limit times according Creep test with 35 Tension test at room temperature Heat treatment to equations 1 and MPa at 750° C. Offset yield Tensile Temperature Time in 2 in minutes Time t_(B) _(—) to Elongation strength strength Elongation in ° C. minutes t_(min) t_(max) break in hours A in % Rp0.2 RM in MPa A 5 in % Hot-worked 511 604 19 I 600 20 27.9 39.2*10⁶ 278 37.6 I 650 20 11.26 31.8*10⁵ 286 35.2 527 622 17 I 700 20 5.00 33.6*10⁴ 260; 264 22.7; 39.1 537 676 13 I 750 20 2.41 44300 263 430.6 514 707 16 I 800 20 1.25 7059 344 25.0 505 699 13 I 850 20 0.69 1328 564 32.2 484 672 17 I 900 20 0.40 289 337 25.3 467 638 18 I 950 20 0.24 71.2 121; 72  31.1; 28.8 451 614 17 R 1000 20 0.15 19.6 10.4 79.5 n.m. n.m. n.m. I 650 240 11.26 31.8*10⁵ 293 36.2 n.m. n.m. n.m. I 700 240 5.00 33.6*10⁴ 233 32.8 n.m. n.m. n.m. I 750 240 2.41 44300 224 23.2 n.m. n.m. n.m. I 800 240 1.25 7059 396 43.2 n.m. n.m. n.m. I 850 240 0.69 1328 181 35.2 n.m. n.m. n.m. I 900 240 0.40 289 45.6 55.7 n.m. n.m. n.m. R 950 240 0.24 71.2 10.8 78.7 n.m. n.m. n.m. I 700 1440 5.00 33.6*10⁴ 645 30.9 n.m. n.m. n.m. For comparison from Example 1 R 1075 20 12.8 69.8 359 494 35 n.m. = not measured

TABLE 4 Limit times according Creep test with 35 Tension test at room temperature Heat treatment to equations 1 and MPa at 750° C. Offset yield Tensile Temperature Time in 2 in minutes Time t_(B) _(—) to Elongation strength strength Elongation in ° C. minutes t_(min) t_(max) break in hours A in % Rp0.2 RM in MPa A 5 in % I 920 28 0.32 162.6 391 38 475 655 18

TABLE 5a Limit times according Creep test with 35 Tension test at room temperature to equations 1 and MPa at 750° C. Offset yield Tensile Batch 161061 2 in minutes Time t_(B) _(—) to Elongation strength strength Elongation Heat treatment t_(min) t_(max) break in hours A in % Rp0.2 RM in MPa A 5 in % I 1050° C./3.4 minutes 0.1 6.0 25.9 50 385 541 27 380 537 28 R 1050° C./20 minutes 0.1 6.0 7.9 83 344 494 31

TABLE 5b Limit times according Creep test with 35 Tension test at room temperature to equations 1 and MPa at 750° C. Offset yield Tensile Batch 161995 2 in minutes Time t_(B) _(—) to Elongation strength strength Elongation Heat treatment t_(min) t_(max) break in hours A in % Rp0.2 RM in MPa A 5 in % I 1050° C./3.4 minutes 0.1 6.0 33.5 89.0 399 538 26 R 1050° C./20 minutes 7.7 92.1 331 475 29

TABLE 6 Limit times according Tension test at room temperature to equations 1 and Creep test at 700° C. Offset yield Tensile Batch 161061 2 in minutes in hours Stress in Time t_(B) _(—) to Elongation strength strength Elongation Heat treatment t_(min) t_(max) MPA break in hours A in % Rp0.2 RM in MPa A 5 in % R 1075° C./22 minutes 40 228 51 367 502 35 R 1075° C./22 minutes 60 8.1 43.1 367 502 35 R 1075° C./22 minutes + 0.083 5601 40 104 72.6 700° C./4 hours R 1075° C./22 minutes + 0.083 5601 60 6.3 63.5 700° C./4 hours

TABLE 7a Limit times according Creep test with 40 Tension test at room temperature Heat treatment to equations 1 and MPa at 750° C. Offset yield Tensile Temperature Time in 2 in minutes Time t_(B) _(—) to Elongation strength strength Elongation in ° C. minutes t_(min) t_(max) break in hours A in % Rp0.2 RM in MPa A 5 in % I 750 20 2.41 44300 128 34.7 514 707 16 I 800 20 1.25 7059 189 27.2 505 699 13 I 850 20 0.69 1328 296 32.1 484 672 17 I 900 20 0.40 289 174 35.4 467 638 18 R 1000 20 0.15 19.6 8.8 78.7 n.m. n.m. n.m. I 750 120 2.41 44300 150 31.7 n.m. n.m. n.m. I 800 120 1.25 7059 227 26.2 n.m. n.m. n.m. I 850 120 0.69 1328 133 29.7 n.m. n.m. n.m. I 900 120 0.40 289 32.7 40 n.m. n.m. n.m. n.m. = not measured

TABLE 7b Limit times according Creep test with 35 Tension test at room temperature Heat treatment to equations 1 and MPa at 750° C. Offset yield Tensile Temperature Time in 2 in minutes Time t_(B) _(—) to Elongation strength strength Elongation in ° C. minutes t_(min) t_(max) break in hours A in % Rp0.2 RM in MPa A 5 in % I 750 240 2.41 44300 182 32.2 n.m. n.m. n.m. I 800 240 1.25 7059 163; 135 27.5; 26.5 n.m. n.m. n.m. I 850 240 0.69 1328 57 33.9 n.m. n.m. n.m. I 750 480 2.41 44300 152 43.4 n.m. n.m. n.m. I 800 480 1.25 7059 169 26.5 n.m. n.m. n.m. I 850 480 0.69 1328 35 33 n.m. n.m. n.m. I 750 960 2.41 44300 139 24.2 n.m. n.m. n.m. I 750 1440 2.41 44300 82 25.5 n.m. n.m. n.m. I 800 1440 1.25 7059 46 46.1 n.m. n.m. n.m. I 750 5760 2.41 44300 54 52.9 n.m. n.m. n.m. I 800 5760 1.25 7059 17.5 50.3 n.m. n.m. n.m. n.m. = not measured

TABLE 8a Limit times according Creep test with 40 Tension test at room temperature Heat treatment to equations 1 and Degree of MPa at 750° C. Offset yield Tensile Temperature Time in 2 in minutes working in Time t_(B) _(—) to Elongation strength strength Elongation in ° C. minutes t_(min) t_(max) % break in hours A in % Rp0.2 RM in MPa A 5 in % I 1050 3.4 0.1 6.0 0.0 33.5 89 399 538 26 R 1050 2.8 0.1 6.0 5.0 10.6 45.8 581 616 16 R 1050 2.8 0.1 6.0 10.0 10.9 68.6 663 685 8 R 1050 2.8 0.1 6.0 20.0 9.4 71.8 725 745 6 R 1050 2.8 0.1 6.0 40.0 12.0 85.9 811 832 4 n.m. = not measured

TABLE 8b Creep test with 40 Tension test at room temperature Degree of Heat treatment MPa at 750° C. Offset yield Tensile working in Temperature Time in Deformation Time t_(B) _(—) to Elongation strength strength Elongation % in ° C. minutes in % break in hours A in % Rp0.2 RM in MPa A 5 in % I 0.0 1075 10 0.0 10.8 50.4 338 479 28 I 2.8 1075 10 2.8 86 23.3 506 545 17 I 2.8 1075 10 2.8 53 19.8 506 545 17 I 5.0 1075 10 5.0 75 18.9 548 571 15 I 5.0 1075 10 5.0 110 49.5 548 571 15 R 10.0 1075 10 10.0 137 31.7 650 555 8 R 20.0 1075 10 20.0 108 23.9 739 750 5 R 20.0 1075 10 20.0 73 60.6 739 750 5 R 40.0 1075 10 40.0 61 24.5 831 837 3 R 40.0 1075 10 40.0 49 34 831 837 3 n.m. = not measured 

1. Method for production of a component, from an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, in that a semifinished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures≧the solution-annealing temperature, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in a second step mechanical working of the semifinished product in the range from 0.05 to 99% is performed and in a subsequent step Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form, by the fact that the component made from the worked semifinished product is brought to an application temperature between 550° C. and 1000° C. by heating at 0.1° C./min to 1000° C./min.
 2. Method for production of a component from an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, in that a semifinished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures≧the solution-annealing temperature, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in a second step mechanical working of the semifinished product in the range from 0.05 to 99% is performed and in a subsequent step Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form by the fact that the worked semifinished product is subjected for a time between t_(min) and t_(max) to a heat treatment in the temperature range between 550° C. and 1060° C. under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800° C. is quenched in the oven, wherein and t_(min) and t_(max) are calculated according to the following formulas: t _(min) =T _(a)·10^((6740/Ta−9.216)) and t _(max) =T _(a)·10^((17960/Ta−15.72)) where T _(a) =T+273.15, and wherein the desired component is made before or after this heat treatment.
 3. Method according to claim 1, wherein, in the first step, the alloy is solution annealed at a temperature ≧1050° C. for longer than 6 minutes.
 4. Method according to claim 1, wherein, in the first step, the alloy is solution annealed at a temperature ≧1060° C. for longer than 1 minute.
 5. Method according to claim 1, wherein semifinished product of the following chemical composition (in % by weight) is thermomechanically treated: Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002-0.1% N 0.002-0.1% S max. 0.01% Fe remainder as well as the usual melting-related impurities, wherein a mechanical deformability at room temperature of >13% is obtained, measured as plastic elongation in the tension test.
 6. Method according to claim 1, wherein only little or even no Fe₂(M, Si) or Fe₇(M, Si)₆ Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are still present in the semifinished product after solution annealing at temperatures the solution-annealing temperature, preferably ≧1050° C. for longer than 6 minutes or ≧1060° C. for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in the initial state before deformation.
 7. Method according to claim 1, wherein the working of the semifinished product takes place by hot working.
 8. Method according to claim 1, wherein the hot working of the semifinished product begins with a starting temperature >1070° C., wherein the last 0.05 to 90% of mechanical deformation is applied between 1000° C. and 500° C.
 9. Method according to claim 1, wherein the hot working of the semifinished product begins with a starting temperature >1070° C., wherein the last 0.05 to 95% of mechanical deformation is applied between 1000° C. and 500° C.
 10. Method according to claim 1, wherein the hot working of the semifinished product begins with a starting temperature >1070° C., wherein the last 0.05 to 99% of mechanical deformation is applied between 1000° C. and 500° C.
 11. Method according to claim 1, wherein the hot working of the semifinished product is followed by cold working.
 12. Method according to claim 1, wherein the working of the semifinished product is carried out by cold working.
 13. Method according to claim 12, wherein the degree of cold working of the semifinished product is 0.05 to 99%.
 14. Method according to claim 12, wherein the cold working of the semifinished product is 0.05 to 95%.
 15. Method according to claim 12, wherein the cold working of the semifinished product is 0.05 to 90%
 16. Method according to claim 1, wherein the mechanical working of the semifinished product is 20 to 99% and then the worked semifinished product is subjected for a time between t_(min) and t_(max) to a heat treatment in the temperature range between 950° C. and 1060° C. under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water and after this the desired component is made with t _(min) =T _(a)·10^((6740/Ta−9.216)) and t _(max) =T _(a)·10^((17960/Ta−15.72)) where T _(a)=T+273.15 and indication of t_(min) and t_(max) in minutes and of heat-treatment temperature T in ° C.
 17. Method according to claim 1, wherein the alloy additionally contains (in % by weight) 0.02 to 0.3% La.
 18. Method according to claim 1, wherein the alloy additionally contains (in % by weight) 0.01 to 0.5% Ti.
 19. Method according to claim 1, wherein the alloy additionally contains 0.02 to 0.3% of one or more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf.
 20. Method according to claim 1, wherein the alloy additionally contains (in % by weight) 0.001 to 0.5% Al.
 21. Method according to claim 1, wherein the alloy additionally contains (in % by weight) 2.0 to 6.0% Al.
 22. Method according to claim 21, wherein the alloy additionally contains (in % by weight) 2.5 to 5.0% Al.
 23. Method according to claim 1, wherein the alloy additionally contains one or more of the elements 0.0001 to 0.07% Mg, 0.0001 to 0.07% Ca, 0.002-0.03% P.
 24. Method according to claim 1, wherein the alloy further contains 0.01 to 3.0% of one or more of the elements Ni, Co or Cu.
 25. Method according to claim 1, wherein the alloy further contains up to 0.005% B.
 26. Method according to claim 1, wherein the iron-chromium alloy, which is thermomechanically treated and which precipitates Laves phases in finely dispersed form, has the following composition containing (in % by weight) Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002-0.03% N 0.002-0.03% S max. 0.005% Fe remainder as well as the usual melting-related impurities.
 27. Method according to claim 1, wherein the alloy additionally contains (in % by weight) 0.02 to 0.2% of the element La.
 28. Method according to claim 1, wherein the alloy additionally contains (in % by weight) 0.02 to 0.2% Ti.
 29. Method according to claim 1, wherein the alloy additionally contains (in % by weight) 0.02 to 0.2% of one or more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf.
 30. Method according to claim 1, wherein the alloy additionally contains (in % by weight) one or more of the elements 0.0001-0.05% Mg, 0.0001-0.03% Ca, 0.002-0.03% P.
 31. Method according to claim 1, wherein the alloy further contains (in % by weight) up to 0.003% B.
 32. Method according to claim 1, wherein (in % by weight) the Nb content is 0.3 to 1.0% and the Si content is 0.15-0.5%.
 33. Method according to claim 1, wherein the W content is replaced entirely or partly by at least one of the elements Mo and/or Ta.
 34. Method according to claim 1, wherein the alloy contains (in % by weight) max. 0.2% V and/or max. 0.005% S.
 35. Method according to claim 1, wherein the alloy contains (in % by weight) max. 0.01% O.
 36. Method according to claim 1, wherein the alloy contains (in % by weight) max. 0.01% of each of the elements Zn, Sn, Pb, Se, Te, Bi and Sb respectively.
 37. Method according to claim 1, wherein the semifinished product is formed by sheet, strip, bar, forging, pipe or wire.
 38. Method according to claim 1, wherein the heat treatment is carried out only after finishing of the component.
 39. Method according to claim 1, wherein, by the thermomechanical treatment of the semifinished product, a particularly high creep strength is produced in the semifinished product and/or in the component with simultaneous elongation >13% in the tension test at room temperature.
 40. Metallic component or semifinished product, consisting of the following chemical composition (in % by weight) Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002-0.1% N 0.002-0.1% S max. 0.01% Fe remainder as well as the usual melting-related impurities, which at the end of a thermomechanical treatment has a deformed microstructure, to the effect that Laves phase(s) is or are embedded in finely dispersed form in the microstructural dislocations of the microstructure, wherein, in a creep test with, in particular, 35 MPa at 750° C. and an elongation of at least 18%, a time to break that exceeds the time to break of a coarse-grained, completely recrystallized microstructure by a factor of at least 1.5 is established in the microstructure.
 41. Metallic component or semifinished product, consisting of the following chemical composition (in % by weight) Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002-0.1% N 0.002-0.1% S max. 0.01% Fe remainder as well as the usual melting-related impurities, which at the end of a thermomechanical treatment has a deformed microstructure, to the effect that Laves phase(s) is or are embedded in finely dispersed form in the microstructural dislocations of the microstructure, wherein, in a creep test with, in particular, 35 MPa at 750° C. and an elongation of at least 18%, a time to break hours that exceeds the time to break of a coarse-grained, completely recrystallized microstructure by a factor of at least 3 is established in the microstructure.
 42. Use of a component produced according to claim 1 as interconnector in a fuel cell.
 43. Use of a component produced according to claim 1 as material in a component, such as a reformer or a heat exchanger or in an ancillary aggregate of a fuel cell.
 44. Use of a component produced according to claim 1 in the exhaust-gas line of a combustion engine.
 45. Use of a component produced according to claim 1 for steam boilers, superheaters, turbines and other parts of a power plant or in the chemical process industry. 